effect os of case car buri zing pm steel

8/10/2019 Effect Os of Case Car Buri Zing Pm Steel

1/19

Effect of Case Carburizing on Mechanical Properties

And Fatigue Endurance Limits of P/M Steels

George Fillari*, Thomas Murphy*, Igor Gabrielov**

*Hoeganaes CorporationCinnaminson, NJ 08077

**Borg Warner Automotive

Livonia, MI 48150

ABSTRACT

Case carburizing has long been a basic technique for the improvement of the wear and fatigue

resistance and of PM steel components. The key to the successful improvement in carburizing,

however, is understanding, and interpreting the microstructure of the carburized case. The main

area for growth in the PM industry is in the high performance gearing applications. The success

of penetrating this area depends upon the ability to understand the key components that effect thefatigue endurance limits of PM materials. This paper will examine and illustrate how tensile

properties and the fatigue endurance limits of a P/M hybrid alloy are affected by alloying

additions and carburizing.

INTRODUCTION

Carburizing is the addition of carbon to the surface of low-carbon steels at temperatures generally

between 850 C and 950 C (1560 F and 1740 F), at which austenite, with its high solubility for

carbon, is the stable crystal structure. Hardening is accomplished when the high-carbon surface

layer is quenched to form martensite so that a high-carbon martensitic case with good wear and

fatigue resistance is superimposed on a tough, low-carbon steel core. [1]

Case depth of carburized steel is a function of carburizing time and the available carbon potential

at the surface. [2] When prolonged carburizing times are used for deep case depths, a high carbon

potential produces a high surface-carbon content, which may thus result in excessive retained

austenite or free carbides. These two microstructural elements both have adverse effects on the

distribution of residual stress in the casehardened part. Consequently, a high carbon potential may

be suitable for short carburizing times but not for prolonged carburizing.

In regards to fatigue properties, Low retained austenite content and fine austenitic grain sizes,

which create a microstructure of finely dispersed retained austenite and tempered martensite,

prevent nucleation of fatigue cracks, or retard fatigue crack initiation until very high stress levels

are reached. In contrast, low-stress applications that fracture at low cycles is related to highretained austenite levels and coarse austenite grain sizes. [1]


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TEST PROGRAM

The test program was intended to investigate the performance levels achievable on tensile

properties and rotating bend fatigue response of a hybrid PM steel by using a secondary heat-

treatments such as case carburizing. This program was divided into three parts.

1. Investigate alloying elements such as graphite and nickel on mechanical properties.

2. To investigate the effect of carburizing times on properties.

3. To estimate the percentage of retained austenite in the carburized case.

EXPERIMENTAL PROCEDURE

The compositions of the test materials that were evaluated in this study are listed in Table I.

The base powder was water atomized and pre-alloyed with 0.85 w/o molybdenum. The 85HP was

premixed with nickel and graphite. The nickel used was INCO 123 and the graphite was Asbury3203H. Each premix contained 0.75 w/o Lonza Acrawax C as the lubricant system.

Table I.Premix Compositions

Premix Base Powder Nickel Graphite Acrawax C

ID Bal. w/o w/o w/o

A Ancorsteel 85HP 2.0 0.15 0.75

B Ancorsteel 85HP 2.0 0.30 0.75

C Ancorsteel 85HP 4.0 0.15 0.75

D Ancorsteel 85HP 4.0 0.30 0.75

TEST SPECIMEN / COMPACTION AND SINTERING

Tensile dog-bone, impact, and fatigue samples were compacted to a density of 7.20 g/cm3.

Green density, sintered density, and transverse rupture strength was determined from the average

of five compacted transverse rupture (TRS) specimens (ASTM B-528). Tensile strength, yield

strength, and maximum elongation were obtained from the average of five dog-bone tensile

samples (ASTM E-8). Apparent hardness measurements were performed on the surface of the

dog-bone tensile samples using a Rockwell hardness tester. All measurements were conducted

using the HRA scale for ease of comparison.

All test pieces were sintered under production conditions in an Abbott continuous belt high

temperature furnace at the Hoeganaes R&D facility, in Cinnaminson, NJ.

The sintering condition used for the test specimen is listed below.


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SINTERING CYCLE

Sintering Temperature: 1150 C (2100 F) (1260 C (2300 F) for premix 3)

Atmosphere: 90 v/o N2- 10 v/o H2Time in Hot Zone: 20 minutes

For the samples that were carburized, the parameters are listed below.

CARBURIZING CYCLE 1

Temperature: 925 C (1700 F) vacuum furnaceTime at Temperature 180 minutes,

Quench: Pressure / Nitrogen

CARBURIZING CYCLE 2

Temperature: 925 C (1700 F)Time at Temperature 240 minutes,

Quench: Pressure / Nitrogen

All samples were tempered at 204 C (400 F) in air for 1hr. prior to testing.

Tensile testing were performed on a 267,000 N (60,000 lb.) Tinius Olsen universal testing

machine with a cross-head speed of 0.635 mm/min (0.025 in/min). Elongation values were

determined by utilizing an extensometer with a range of 0 - 20%. The extensometer was attached

to the samples up to failure.

Rotating bending fatigue samples were pressed to a density of 7.20 g/cm3, and machined from

blanks that were sintered at 1150 C (2100 F) under an atmosphere of 90 v/o N2-10 v/o H2.

The heat treated fatigue samples were rough machined following sintering then heat treated,

finished ground, and polished to size. The dimensions of the specimen used for this analysis,

along with allowable dimensional tolerances, are shown in Figure 1.

Fatigue testing was performed on six randomly selected Fatigue Dynamics RBF-200 machines ata rotational speed of 8000 rpm. These rotating bending machines are of the mechanical and non-

resonant type and are an efficient means of inducing fatigue in a specimen of round cross section.

[3]. A staircase method was used utilizing 30 samples and a run-out limit of 107cycles. The

staircase method of testing was regulated so that there were both failures and run-outs at a

minimum of two stress levels. [4] The percentage of failures for each stress level was calculated

and plotted on a log-normal graph. From these plots, the fatigue endurance limit (FEL) at 50%

and 90% was determined by linear extrapolation. The 50% FEL represents the stress level where

50% of the specimens will break and 50% will run-out. The 90% FEL represents the stress level

where 90% of the specimens will run-out and 10% will break.


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Figure 1. Dimensions of Rotating Bending Fatigue Specimen

RESULTS AND DISCUSSION

The mechanical properties of the alloys evaluated are summarized in Tables II trough Table VI.

Shown in Figure 2 and 3 are the effects of alloy content and carburizing times on ultimate tensile

and yield strengths for the alloys tested.

In the as sintered condition, the tensile strengths for the alloys at a sintered density of 7.20 g/cm3

are 490 MPa (71*103psi), 559 MPa (81*103psi), 621 MPa (90*103psi) and 738 MPa (107*103

psi) for alloys A, B, C and D. Elongations are in the range of 3.1 5.0 %. For the samples that

were carburized in the first cycle, the tensile strengths were increases between 23 70 % to 835

MPa (121*103psi), 850 MPa (123*103psi), 850 MPa (123*103psi) and 910 MPa (131*103psi),

when compared to the as sintered condition. Samples that were subjected to the second

carburizing cycle resulted in increases in tensile strengths between 5 20 % to 1000MPa(145*103psi), 891 MPa (129*103psi), 987 MPa (143*103psi) and 959 MPa (139*103psi) when

compared to the first cycle. Increases in yield strengths also follow the same trend. High

temperature sintering for alloy 3 resulted in an increase in tensile strength from 3 13 %.

Table II.As Sintered Properties for The Alloys Tested at 1150 C (2100 F).

Sintered 0.002

Condition Material Density HRA UTS OFFSET Elong

ID (g/cm) (MPa/103psi) (MPa/10

3psi) %

1 7.24 47 490/71 352/51 4.6

2 7.22 50 559/81 386/56 4.33 7.25 51 621/90 407/59 5.0

4 7.24 54 738/107 524/76 3.1

As Sintered


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Table III.Tensile Properties for Samples Subjected to Both Carburizing Cycles. The Alloys

Tested. Sintered at 1150 C (2100 F)

Sintered Apparent 0.20%

Condition Material Density Hardness UTS Offset Elong

ID (g/cm3) HRA (MPa/10

3psi) (MPa/10

3psi) %

A 7.22 73 835/121 731/106 0.8B 7.22 74 850/123 842/122 0.8

C 7.26 70 850/123 607/88 1.1

D 7.27 71 910/131 607/88 1.2

A 7.22 73 1000/145 918/133 1.0

B 7.21 73 890/129 752/109 0.9

C 7.27 73 987/143 718/104 1.1

D 7.26 72 959/139 676/98 1.1

Cycle 1Tempered

@ 205 C

(400 F)Cycle 2

Tempered

@ 205 C

(400 F)

Table IV.Alloy C Tensile Properties for Samples Subjected to Both Carburizing Cycles.

Sintered at 1260 C (2300 F)

Carburizing Sintered Apparent 0.20%Condition & Density Hardness UTS Offset Elong

Tempering Temp (g/cm3) HRA (MPa/10

3psi) (MPa/10

3psi) %

Cycle 1 @ 205 C 7.31 73 966/140 621/90 1.3

Cycle 2 @ 205 C 7.30 73 1021/148 711/103 1.1

300

400

500

600

700

800

900

1000

1100

Alloy A Alloy B Alloy C Alloy D

UTS(MPa)

AS SINT CYCLE 1 CYCLE 2

X

X

Figure 2. Ultimate tensile strength as a function of alloy content and carburizing cycle. Samples

sintered at 1150 C (2100 F) - X indicates the strength of alloy sintered at 1260 C (2300 F)


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200

300

400

500

600

700

800

900

1000

Alloy A Alloy B Alloy C Alloy D

YieldStrength(MPa)

AS SINT CYCLE 1 CYCLE 2

X

X

Figure 3. Yield strength as a function of alloy content and carburizing cycles. Samples sintered at1150 C (2100 F) - X indicates the strength of alloy sintered at 1260 C (2300 F)

Rotating bending fatigue data were collected on samples compacted to a density of 7.20 g/cm3.

The fatigue endurance limits determined for the materials, along with sintered densities are shown

in Table V. Figure 4, illustrates the fatigue performance for the alloys tested along with the

fatigue ratios for 90%. What is interesting to note is that the carburized samples resulted in

fatigue endurance limits equivalent to or greater than samples that were machined from AISI

8620 wrought. The wrought samples were machined in the principal working direction

(longitudinal) and perpendicular to the principal working direction (transverse). [6]

Table V. Rotating Bending Fatigue Results sintered at 1150 C (2100 F) Tempered at 205 C

(400 F)

Carburizing Sintered

Condition Material Density Fatigue Ratio

ID (g/cm3) 50% 90% 90% Survival

A 7.22 462/67 449/65 0.53

B 7.22 469/68 455/66 0.55

C 7.26 497/72 483/70 0.58

D 7.27 504/73 490/71 0.53

C* 7.31 504/73 497/72 0.51

A 7.22 414/60 393/57 0.39

B 7.21 455/66 462/67 0.52

C 7.27 483/70 469/68 0.47

D 7.26 435/63 414/60 0.43

C* 7.30 490/71 476/69 0.46

Survival Limits

Cycle 1

Cycle 2

(* Alloy C Sintered at 1260 C (2300 F)


7/19

300

350

400

450

500

550

600

800 900 1000 1100 1200 1300 1400

Tensile Strength (MPa)

FatigueEnduranceLimit(MPa)

A-1 (.53 UTS)

B-1 (.55 UTS)

C-1 (.58 UTS)

D-1 (.53 UTS)

B-2 (.52 UTS)

C-1 @ 1260 C (.51 UTS)

C-2 (.47 UTS)

C-2 @ 1260 C (.46 UTS)

D-2 (.43 UTS)

A-2 (.52 UTS)

Wrought

Long. .35 UTS)

Wrought

Trans. .26 UTS)

Figure 4.The relationship of fatigue endurance limits and tensile strength of the alloys tested,

along with the fatigue ratios for 90% compared to AISI 8620 wrought.

The etched microstructure in the as-sintered condition are illustrated in Figure 5(a-d). Figure 5a.

consist mostly of divorced pearlite, ferrite with some martensitic regions. Figure 5b is similar

with nickel rich regions and bainite and martensite. Figure 5c is similar to Figure a and b. The

microstructure of Figure 5d consist of divorced pearlite, areas of martensite, and some bainite.

Figures 6 through 10 compares the microstructure of the samples subjected to both carburizing

cycles. These microstructures are taken from the center of the fatigue samples at the reduced

sections. From Figures 6 10 (a-b) you can clearly see the carburized case that resulted from the

first cycle. These samples resulted in a case depth of about 1100 m (Alloy A), 900 m (Alloy

B), 1000 m (Alloy C), 1100 m (Alloy D), and 1000 m (Alloy C sintered at 1260 C (2300 F),

respectively. When compared to the samples subjected to the second carburizing cycle (Figures 6

10 (e-f), these samples were through hardened with no evidence of a definitive carburized case.

The etched structure for the alloys carburized in the first cycle are similar. In the reduced section

(Figure 6 10 c,) consist of acicular martensite, retained austenite with some nickel rich regions.

The etched structure of the softer core of alloy A (Figure 6d), consist of lathe martensite, bainite

and nickel rich regions. The core of alloy B and C (Figure 7d and 8d) consist of divorced pearlite

and nickel rich regions. The core of alloy D (Figure 9d) consist of mostly lathe marensite, nickelrich regions, bainite, and unresolved pearlite. The core of alloy C sintered at 1260 C (2300 F),

(Figure 10d) consist of lathe martensite, nickel rich regions and divorced pearlite.

The microstructures for the alloys that were through hardened in the second cycle are also similar.

In the reduced sections (Figure 6-10, g-h) consist mostly of acicular martensite, retained austenite

with nickel rich regions.


8/19

Figure 11 illustrates the Vickers Hardness HV of the alloys tested. Hardness readings were taken

from the surface of the case to the core in increments of 0.0002 in. As a result of the through

hardening, the samples subjected to the second cycle resulted in a higher hardness (HV) though

out the sample. The samples from the first cycle initially were harder at the surface and gradually

decreased in hardness as the distance increased from the surface to the core.

In terms of the effect of the nickel and graphite. In the as-sintered condition, ultimate tensilestrength increased with nickel content. For the alloys that contained 0.15 w/o graphite, as the

nickel content increase from 2.0 to 4.0 w/o the UTS increase more than 25 % to 621 MPa (90*103psi). For the alloys that contained 0.30 w/o graphite, as the nickel content was increased to 4.0

w/o the UTS increased more than 30 % to 738 MPa (107*103psi). The same trends also resulted

for the yield strength. For the alloys at 0.15 w/o graphite, the yield increased more than 15 % to

407 MPa (59*103psi). At 0.30 w/o graphite, yield strength increased more than 35 % to 524 MPa

(76*103psi).

For the samples that were carburized (cycle 1) the increases were minimal. For the alloys that

contained 0.15 w/o graphite, as the nickel content increase from 2.0 to 4.0 w/o the UTS increase

were about 2 % to 850 MPa (123*103psi). For the alloys that contained 0.30 w/o graphite, as the

nickel content was increased to 4.0 w/o the UTS increased about 5 % to 910 MPa (131*103psi).The yield strength on the other hand decreased as the nickel content was increased to 4.0 w/o. For

the alloys at 0.15 w/o graphite, the yield strength decreased more than 15 % to 607 MPa (88*103

psi). At the 0.30 w/o graphite contents, yield strength decreased more than 25 % to 607 MPa

(88*103psi).

For the samples that were through hardened (cycle 2), at the 0.15 w/o graphite, as the nickel

content increase to 4.0 w/o the UTS decreased to 987 MPa (143*103psi), but was negligible. At

0.30 w/o graphite, as the nickel content was increased to 4.0 w/o the UTS resulted in a slight

increase to 959 MPa (139*103psi). The same trends also resulted for the yield strength as in the

first cycle. For the alloys at 0.15 w/o graphite, the yield decreased more than 20 % to 718 MPa

(104*103psi). At the 0.30 w/o graphite contents, yield strength decreased 10 % to 676 MPa

(98*103

psi).

For the high temperature sintering of alloy C for the first carburizing cycle, the UTS increases

were about 15% to 966 MPa (140*103psi), changes in yield were negligible. For the second

carburizing cycle, a slight increase in UTS to 1021 MPa (148*103psi), changes in yield were

negligible.


9/19

40 m

(a)

40 m

(b)

40 m

(d)

40 m

(c)

Figure 5. Microstructure of the as sintered samples: (a)Alloy A, (b) Alloy B, (c) Alloy C, (d)

Alloy D. Sintered at 1120 C (2050 F)


10/19

500 m

(e)

500 m

(a)

900 m

(b)

900 m

(f)

40 m

(g)

40 m

(c)

40 m

(d)

40 m

(h)

Figure 6. Microstructures of alloy A: (a-c (case) d (core) subjected to carburizing cycle-1) (e-g(case) h (core) subjected to carburizing cycle-2) Sintered at 1150 C (2100 F)


11/19

500 m

(a)

900 m

500 m

(f)

(e)

900 m

(b)

40 m

(c)

40 m

(g)

40 m

(d)

40 m

(h)

Figure 7. Microstructures of alloy B: (a-c (case) d (core) subjected to carburizing cycle-1) (e-g

(case) h (core) subjected to carburizing cycle-2) Sintered at 1150 C (2100 F)


12/19

500 m

(a)

500 m

(e)

900 m

(b)

900 m

(f)

40 m

(g)

40 m

(h)

40 m

(c)

40 m

(d)

Figure 8. Microstructures of alloy C: (a-c (case) d (core) subjected to carburizing cycle-1) (e-g(case) h (core) subjected to carburizing cycle-2) Sintered at 1150 C (2100 F)


13/19

500 m

(a)

500 m

(e)

900 m

(f)

40 m

(g)

900 m

(b)

40 m

(c)

40 m

(d)

40 m

(h)

Figure 9. Microstructures of alloy D: (a-c (case) d (core) subjected to carburizing cycle-1) (e-g



14/19

500 m

(a)

500 m

(e)

900 m

(b)

900 m

(f)

40 m

(g)

40 m

(h)

40 m

(c)

40 m

(d)

Figure 10. Microstructures of alloy C: (a-c (case) d (core) subjected to carburizing cycle-1) (e-g



15/19

0

100

200

300

400

500

600

700

800

0.0002 0.0006 0.001 0.0014 0.0018

Distance from Surface (in.)

Vick

ersHardness

Cycle 1

Cycle 2

a

0

100

200

300

400

500

600

700

800

0.0002 0.0006 0.001 0.0014 0.0018


Vick

ersHardness

Cycle 1

Cycle 2

(b)

0

100

200

300

400

500

600

700

800

0.0002 0.0006 0.001 0.0014 0.0018


Vicker

sHardness

Cycle 1

Cycle 2 d0

100

200

300

400

500

600

700

800

0.0002 0.0006 0.001 0.0014 0.0018


VickersHardness

Cycle 1

Cycle 2

Cycle 1 @ 1260 C

Cycle 2 @ 1260 C

c

Figure 11. Vickers Hardness (HV) of the samples evaluated. Sintered at 1150 C (2100 F) a)

Alloy A, b) Alloy B, c) Alloy C 1260 C (2300 F) d) Alloy D

ESTIMATE OF RETAINED AUSTENITE CONTENT IN CARBURIZED SAMPLES

A study was conducted to evaluate the effect of surface carburization of test bars for the alloys

tested. The metallographic test pieces used in this exercise were cross-sections cut from sintered

and carburized impact bars.

Cross-sections from the carburized impact bars from the first cycle were cut, mounted, ground

and polished using well-established metallographic practices. The prepared surfaces were then

etched using a combination of 2 v/o nital and 4 v/o picral. Preliminary visual examination

confirmed the presence of the carburized case where the microstructure consisted of martensite

and retained austenite, nickel rich regions, and a few small areas of bainite. The etched surfaces

were further prepared by stain etching with an aqueous solution of 25 w/o sodium bisulphite. A

contrast was created in the microstructure between the martensite, which was colored by the stain

etch, and the featureless white retained austenite/Ni-rich regions.

An automated image analysis system was used to measure the amount of the unstained white

areas (retained austenite). During the analysis, the features coinciding with a near white detection

setting were separated into large and small feature groups. The small features were < 30 m2and

appear as small, angular needles located between the martensite needles. The large features, > 30

m2and appear to be a combination of Ni-rich regions and areas containing diffused Ni with the

elevated carbon content from the carburization process.

2and appear to be a combination of Ni-rich regions and areas containing diffused Ni with the

elevated carbon content from the carburization process.


16/19

Area measurements were made in strips of fields along the edge of each cross-section. Each strip

was a single field in height and 30 fields across. The first strip was positioned beneath the sample

edge at a distance of 150 m from the edge. The second strip was extended 150 m 300 m

into the sample. Each strip was approximately 6 mm long. In making the measurements, the total

volume of the compact was considered, including porosity. It was not subtracted from the

calculations. Table 6 list the volume percent measurements of the two white feature sizes

(retained austenite) and the averages at specific distances within the part along with the averageminus the volume of the Ni-regions.

Table VI. Area Percent Measurements

Retained Aust

+ Ni-rich Depth Average w/o

Sample-Strip Feature Size Average (v/o) micron Ret. Aust.

1 - 1 < 30 micron 3.24

> 30 micron 7.61

1 - 2 < 30 micron 1.99

> 30 micron 6.36

2 - 1 < 30 micron 3.73

> 30 micron 11.39

2 - 2 < 30 micron 2.44> 30 micron 7.25

3 - 1 < 30 micron 2.88

> 30 micron 26.99

3 - 2 < 30 micron 2.50

> 30 micron 21.99

4 - 1 < 30 micron 2.45

> 30 micron 22.14

4 - 2 < 30 micron 2.01

> 30 micron 19.26300

300

150

300

150

150

300

150

18.17

12.79

12.89

9.57

5.75

3.25

10.02

4.59

The averages of the small fractions remain consistent throughout the examination, staying

between 2 and 4 v/o for all samples. The major difference appeared in the measurement of the

large features, especially in the 4 w/o Ni premixes. An interpretation problem was seen in thesetests where the separation of the retained austenite from the areas containing very high Ni

contents could not be made because of any clear distinction between chemical and/or

microstructural composition was apparent. The increased alloy content from the diffused Ni

coupled with the increase in carbon content from the carburization treatment probably caused the

Ms temperature in some regions to be lowered sufficiently to prevent transformation to

martensite.

Further tests were preformed to determine the amount of Ni-rich regions present in the premixes

with both Ni contents and in the areas unaffected by the carburization treatment. Core

microstructures for premixes 1 (2w/o Ni) and 3 (4w/o Ni) were used as non-carburized examples.It was thought the possibility of comparing and subtracting the amount of core Ni-rich phases

from the total may help quantify the higher alloyed retained austenite in the carburized case.

To accomplish this, the samples were prepared and re-etched with the nital/picral combination.

The etch/stain procedure described previously to develop the microstructural contrast wasineffective on the core microstructures because the constituents within the cores were obviously

different from the carburized cases. The cores consist of ferrite, divorced pearlite, bainite, some

martensite, Ni-rich regions, and probably retained austenite surrounding the Ni-rich regions.


17/19


18/19

3. The 90 % fatigue survival limits decreased for the 2ndcarburizing cycle. The limits variedbetween 39 to 52%. High temperature sintering reduced the 90% fatigue limits to 46%. With

respect to tensile properties, high temperature sintering for the 2ndcarburizing cycle resulted

in a negligible increase in ultimate and yield strength.

4. For both carburizing cycles, increasing the nickel content from 2.0 w/o to 4.0 w/oresulted in minimal if any increases in tensile and fatigue properties.

5. The carburized samples resulted in fatigue endurance limits equivalent to or greater thansamples that were machined from AISI 8620 wrought.

6. Nickel plays an important role in the formation of retained austenite. The samples thatcontained 4.0 w/o Ni resulted in an increase in retained austenite for more than 70 % to for

the samples that contained 0.15 w/o Gr and more than 50 % for the sample that contained

0.15 w/o Gr (in the range of 300 m).At this level of retained austenite, between 5 15%, any differences in mechanical properties appear to be negligible.

ACKNOWLEGEMENTS

The authors wish to acknowledge the contributions of Gerald Golin and Thomas Murphy to this

paper and their timely preparation of the metallographic samples is much appreciated.


19/19

REFERENCES

1. Krauss, G., Principles of Heat Treatment of Steels, American Society for Metals, Vol. 1, pp.251

2. Case Hardening of Steel, ASM International, Vol. 1, pp. 4, 1987

3. Manual on Fatigue Testing, University Microfilms, Inc., Baltimore, MD, 1949.

4. Rice, R.C., Fatigue Data Analysis, ASM International, Metals Handbook, Vol. 8, 9thEdition, pp. 695-720, 1985.

5. W. Jandeska, R. Slattery, H. Fran, A. Rawlings, P. King, Rolling Contact Fatigue Evaluationof Powder Forged FLN2-4405, To be presented 2005 International Conference on Powdered

Metallurgy and Particulate Materials, Montreal, Canada.

effect os of case car buri zing pm steel

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